This invention relates to the heat treatment of metals and has particular reference to the heat treatment of titanium near alpha alloys.
The search for improved mechanical properties in titanium alloys has normally taken the route of modifying the composition of the alloy to improve the balance of properties available. Titanium alloys have been in existence commercially for a little over 30 years and it is becoming increasingly difficult to design new titanium alloys with improved properties.
Initial improvements were made quite rapidly, but the rate of development has slowed down as the law of diminishing returns takes effect. Undoubtedly improvements will occur in the future. However, even small improvements in properties are valuable in that they enable aero engines to be designed so as to be lighter and hence more fuel efficient. The need for fuel efficiency in aero engines is so great that aero engine designers are looking to use titanium alloys in ever hotter regions of the engine to enable weight savings to be obtained. There is, therefore, a great deal of pressure on the metallurgist to improve the balance of metallurgical properties present in the alloy.
As mentioned above most of the emphasis on improving properties has gone towards modifying the composition of the alloy. Little practical evaluation has been given to modifications to the heat treatment to be used on the alloys. This invention is, however, concerned with the improvement in titanium alloys by modifying the heat treatment given to them during their processing.
As in the case of many metals titanium exists mainly in two distinct phases, a so-called alpha phase and a so-called beta phase. The beta phase is more stable at elevated temperatures and the proportions of alpha and beta in various titanium alloys are defined by the composition and heat treatment of the alloys. Certain alloying elements used in titanium stabilise the alpha phase and these are frequently referred to as alpha stabilisers. Other alloying elements stabilise the beta phase and these are frequently referred to as beta stabilisers. Certain titanium alloys consist almost completely of alpha titanium when in equilibrium at room temperature with a trace of beta--less than 5% beta. These alloys are sometimes referred to as near alpha alloys and certain of the alloys are properly regarded as weldable. A near alpha titanium alloy may also be regarded as one containing not more than about 2% by weight of beta stabilisers such as molybdenum copper silicon etc. A more complete definition of a near alpha titanium alloy is an alpha stabilised alloy, that is an alloy containing alpha stabilising elements, with an amount of beta stabiliser which gives a small volume fraction (less than about 5%) of retained beta and which can be beta processed and/or beta heat treated and give acceptable ductility and fracture resistance.
The term "weldable" as used herein is not intended merely to refer to the ability of the metal to be welded directly to itself but is intended to refer to the metal being useable in an aircraft engine in the welded condition. The only two weldable near alpha beta heat treated alloys in existence at the present time are the alloys known as IMI 685, namely the alloy 6% aluminium, 5% zirconium, 0.5% molybdenum, 0.25% silicon, balance titanium and 5331S, namely the alloy 5.5% aluminium, 3.5% tin, 3% zirconium, 1% niobium, 0.25% molybdenum, 0.3% silicon, balance titanium. All percentages as used herein are weight percentages. The near alpha alloys are conventionally used in the solution treated and stress relieved condition. The solution treatment of the alloy 5331S conventionally comprises a treatment at 1050.degree. C. for a time depending on section size--one hour per 2.5 cm. The alloy is then oil quenched and is given a stress relief treatment for two hours at 625.degree. C. although the exact stress relief time may vary with section. The solution treatment modifies the metallurgical structure of the alloy and the stress relieving treatment stress relieves the alloy from the stresses built up in the alloy during the quenching phase.
It will be appreciated that different types of titanium alloys have different types of heat treatment. Thus a conventional heat treatment for a near alpha alloy has been solution treatment in the beta field followed by a stress relieving treatment at a temperature typically in the region 525.degree.-625.degree. C. for a time of about 24 hours. By comparison, however, other types of titanium alloys are given a very different type of heat treatment. Thus an age hardenable titanium alloy, such as titanium plus 21/2% copper, would be given an alpha solution treatment at about 800.degree. C. followed by a nucleation treatment at 400.degree. C. for 8 hours to nucleate the typical "Duralumin" type precipitate and then a further heat treatment at 475.degree. C. for 8 hours to grow the precipitate. The alloy titanium plus 21/2% copper is one which contains only beta stabilisers and is normally treated in the alpha plus beta or alpha plus compound regions of the phase diagram. In effect alloys of this precipitation hardening type rely on forming, at room temperature, a supersaturated solution of copper in the alpha phase. Subsequently the age hardening heat treatments result in the diffusion of copper to precipitation sites and then further precipitation on these sites during subsequent heat treatment.
There are believed to be no commercially used fully beta stable titanium alloys. Experimental alloys such as titanium plus 20% molybdenum plus 10% vanadium are fully beta stabilised. The only heat treatment given to such alloys is to beta solution heat treat. No further heat treatment is given.
A typical metastable beta titanium alloy, such as titanium plus 15% molybdenum would be given a beta solution treatment at a temperature above 25.degree. C. above the beta transus, i.e. 815.degree. C. for the Ti+15% Mo alloy and it would then be water quenched to room temperature. The alloy would then be composed of 100% beta phase. It would then be given a single or duplex ageing to precipitate out from the beta phase either an omega phase or an alpha phase.
Alpha plus beta titanium, such as the alloy titanium plus 6% aluminium plus 4% vanadium is typically heat treated in one of two ways. In one way, the alloy is annealed at a temperature low in the alpha plus beta phase field--i.e. 700.degree. C. to give equiaxed alpha plus retained beta. In the other heat treatment the alloy is solution treated in the alpha plus beta field, air cooled to room temperature and then stress relieved at a single temperature in the range 500.degree. C. to 700.degree. C. to give an equiaxed alpha plus transformed beta structure.
Plain alpha titanium such as commercial purity titanium is simply stress relieved with a single heat treatment in the range 600.degree. C. to 700.degree. C. to given an equiaxed primary alpha structure.
However, it is not possible to equate the heat treatment used for one type of alloy, such as an age hardenable alloy of the titanium plus 21/2 copper type, with that required for another type of alloy, such as a metastable beta or near alpha alloy.
Although practical heat treatments have been developed for near alpha alloys and have been shown to work well it is not certain what is happening in the near alpha alloy when it is heat treated. During the solution treatment it is clear that the alloy is converted into the beta phase and during cooling converts mainly to the alpha phase. However the heat treatment given to stress relieve the alloy after cooling gives rise to numerous types of reactions within the alloy itself.
Thus during the stress relieving process it is quite probable that some form of ordering is taking place within the alpha matrix and furthermore some amount of precipitation of very fine particles of material is taking place within the matrix. Once precipitated the morphology of the precipitate is altered as the heat treatment persists. Furthermore subcells are formed within the alloy. In addition to changes in relation to the precipitate there are also changes in the composition of the matrix.
The relative speeds of the various reactions alter as the temperature of heat treatment changes and furthermore vary with the time at a given temperature. This makes the prediction of the outcome of a variation in heat treatment very difficult when it is considered on a detailed and practical scale.
A subcell of the type referred to above is basically a subgrain in which there is a small difference in the angle of the atomic planes between one cell and another of the order of 5.degree., whereas for a true grain boundary the angular differences between the atomic planes would normally be 30.degree. or more. A subcell may be regarded as a sign of partial recovery within the alloy caused by small movements of dislocations in the alloy. As the amount of precipitate and the morphology of the precipitate changes, the ability of the precipitate to lock up dislocations also changes, and this again gives rise to variations in the properties of the material.
An important part of the stress relieving treatment given to near alpha alloys is to stress relieve the internal stresses built up in the alloy during the quenching from the solution treatment temperature. These stresses are conventionally relieved by the movement of dislocations within the material and by the reformation of grain boundaries, and consequently the effect of the type of precipitate and its morpholoy on stress relief is a further complication.
Although extending the time of the stress relief treatment or increasing the temperature of the heat treatment reduces the amount of internal stress, it has been found that in near alpha alloys this reduces the creep strength of the alloy very considerably. Thus from Table I it can be seen increasing the temperature of the stress relief treatment from 500.degree. C. to 600.degree. C. whilst keeping the duration of the treatment constant at 24 hours led to a doubling of the creep extension a marginal fall in the strength of the alloy and a significant reduction in ductility. The alloy being tested was the near alpha alloy IMI 685. All the material was solution treated at 1050.degree. C. and oil quenched.
TABLE I __________________________________________________________________________ Stress Relief Creep T.P.S., 520.degree. C. 310 N .multidot. mm.sup.-2 0.2% PS UTS EL5D R in A Treatment 100 hrs, % Nmm.sup.-2 Nmm.sup.-2 % % __________________________________________________________________________ 24 hrs/500.degree. C. -- 888 1016 12 23 " 0.063 *922 1013 11.5 19* 24 hrs/575.degree. C. -- 900 1013 7 16 " 0.119 *936 1018 5 7* 24 hrs/600.degree. C. -- 883 999 8 13 " 0.124 *939 1008 3 8* __________________________________________________________________________ *All post creep tensile test samples had their surfaces retained.
For each heat treatment pair, the upper line refers to material which has not been creep tested the lower line for material which has been creep tested.
The same effect of a fall in the creep strength was observed when the time of the stress relief treatment was increased at constant temperature.
Table II, below shows that increasing the stress relief time at a constant temperature gives an increase in strength but a marked reduction in creep strength. The alloy tested was 5331S which had been solution treated at 1050.degree. C. for 2 hours and then oil quenched.
TABLE II __________________________________________________________________________ Creep T.P.S., Stress Relief 540.degree. C./300 Nmm.sup.-2 0.1% PS 0.2% PS UTS EL5D R in A Heat Treatment 100 hr % 300 hr % Nmm.sup.-2 Nmm.sup.-2 Nmm.sup.-2 % % __________________________________________________________________________ 2 hrs/625.degree. C. -- -- 845 865 999 14 17.5 " 0.084 0.256 *913 932 1027 7.5 10* 4 hrs/625.degree. C. -- -- 843 867 995 12 14 " 0.135 0.305 *917 937 1030 8.5 8.5* 8 hrs/625.degree. C. -- -- 861 881 1001 11 16 " 0.164 0.351 *926 945 1038 6 7* __________________________________________________________________________ *All post creep tensile test samples had their surfaces retained.
It will be appreciated that an alloy which has a good creep resistance is one which will extend as little as possible under creep loading conditions, i.e. the value of creep T.P.S. (total plastic strain) should be as low as possible.
It has now been discovered that the properties of near alpha alloys, and in particular 5331S, can be improved by modifying the heat treatment given heretofor to alloys of this type. In particular it has been found that the strength and creep resistance of the alloy can be improved by modification to the known heat treatment.